Nanoscale ion storage materials including co-existing phases or solid solutions

ABSTRACT

Nanoscale ion storage materials are provided that exhibit unique properties measurably distinct from their larger scale counterparts. For example, the nanoscale materials can exhibit increased electronic conductivity, improved electromechanical stability, increased rate of intercalation, and/or an extended range of solid solution. Useful nanoscale materials include alkaline transition metal phosphates, such as LiMPO 4 , where M is one or more transition metals. The nanoscale ion storage materials are useful for producing devices such as high energy and high power storage batteries, battery-capacitor hybrid devices, and high rate electrochromic devices.

RELATED APPLICATIONS

This application is a continuation of U.S. patent application Ser. No.11/396,515 which claims the benefit of priority to U.S. Application No.60/706,273, filed Aug. 8, 2005, and U.S. Application No. 60/741,606,filed Dec. 2, 2005, the contents of which are incorporated herein byreference.

BACKGROUND

1. Field

The field includes ion storage materials, and in particular nanoscaleion storage materials useful in devices such as batteries.

2. Summary of Related Art

Ion storage materials are widely employed in storage batteries and otherelectrochemical devices. Various ion storage materials are known,including alkaline transition metal phosphates. This class of compoundstypically has crystal specific gravity values of about 3 g/cm³ to about5 g/cm³, and can crystallize in a number of structure types. Examplesinclude ordered or partially disordered structures of the olivine(A_(x)MXO₄), NASICON (A_(x)(M′,M″)₂(XO₄)₃), VOPO₄, LiVPO₄F, LiFe(P₂O₇)or Fe₄(P₂O₇)₃ structure types, wherein A is an alkali ion, and M, M′ andM″ are metals. Many such compounds have relatively low electronicconductivity and alkali ion conductivity, which are less than ideal forelectrochemical applications. Many such compounds also exhibit limitedsolid solution range. For example, LiFePO₄ has been widely reported inthe scientific literature to have an extremely limited range of solidsolution at room temperature.

“Nanocrystalline” ion storage materials have been reported in theliterature. For example, Prosini et al. in “A New Synthetic Route forPreparing LiFePO₄ with Enhanced Electrochemical Performance,” J.Electrochem. Soc., 149:A886-A890 (2002), describe LiFePO₄ of 8.95 m²/gspecific surface area as nanocrystalline. However, these materials,while somewhat improved, have not been of sufficiently small size scaleto provide substantially different properties (e.g., near-theoreticalcapacity at high rates in excess of 5C) compared to their larger scalecounterpart conventional ion storage materials.

SUMMARY

Nanoscale ion storage materials are provided that exhibit uniqueproperties measurably distinct from their larger scale counterparts. Forexample, the disclosed nanoscale materials can exhibit increasedelectronic conductivity, improved electromechanical stability, increasedrate of intercalation, and an extended range of solid solution.

In one aspect, a lithium transition metal phosphate material for use asan ion storage material is provided, including at least two co-existingphases, including a lithium-rich transition metal phosphate phase and alithium-poor transition metal phosphate phase, wherein the percentagemolar volume difference between the two phases is less than about 6.5%.

In one or more embodiments, the percentage molar volume differencebetween the two phases of the lithium transition metal phosphatematerial is less than about 6.40%, or less than about 6.25%, or lessthan about 5.75%, or less than about 5.5%.

In one or more embodiments, the at least two existing phases of thelithium transition metal phosphate material are crystalline and aredefined by a unit cell having lattice parameters for each principalaxis, and wherein the difference in lattice parameters for at least twoprincipal axes of the unit cells are less than 3%.

In one or more embodiments, the difference in lattice parameters for allprincipal axes of the unit cells are less than 4.7%, or the differencein lattice parameters for all principal axes of the unit cells are lessthan 4.5%, or the difference in lattice parameters for all principalaxes of the unit cells are less than 4.0%, or the difference in latticeparameters for all principal axes of the unit cells are less than 3.5%.

In one or more embodiments, the difference in the smallest product oflattice parameters for any two principal axes of lithium transitionmetal phosphate material is less than 1.6%, or the difference in thesmallest product of lattice parameters for any two principal axes isless than 1.55%, or the difference in the smallest product of latticeparameters for any two principal axes is less than 1.5%, or thedifference in the smallest product of lattice parameters for any twoprincipal axes is less than 1.35%, or the difference in the smallestproduct of lattice parameters for any two principal axes is less than1.2%, or the difference in the smallest product of lattice parametersfor any two principal axes is less than 1.0%.

In one or more embodiments, the difference in the largest product oflattice parameters for any two principal axes of lithium transitionmetal phosphate material is greater than 4.7%, or the difference in thelargest product of lattice parameters for any two principal axes isgreater than 4.8%, or the difference in the largest product of latticeparameters for any two principal axes is greater than 4.85%.

According to one embodiment, the nanoscale materials have a plane formedby any of the principal axes of the crystal along which the strainmeasured as a change in the area is less than about 1.6%, or less thanabout 1.5%, or less than about 1.4%. According to another embodiment,none of the planes formed by any of the principal axes of the crystalhave such a strain exceeding 8%, or 7.5%, or 6%.

In one or more embodiments, the lithium transition metal phosphatematerial has a specific surface area of at least about 20 m²/g, or atleast about 35 m²/g, or at least about 50 m²/g.

In one or more embodiments, the lithium transition metal phosphatematerial is selected from the group consisting of ordered or partiallydisordered structures of the olivine (A_(x)MPO₄), NASICON(A_(x)(M′,M″)₂(PO₄)₃), VOPO₄, LiVPO₄F, LiFe(P₂O₇) or Fe₄(P₂O₇)₃structure types, wherein A is an alkali ion, and M, M′ and M″ aretransition metals.

In one or more embodiments, the lithium transition metal phosphatematerial has an overall composition of Li_(1-x)MPO₄, where M comprisesat least one first row transition metal selected from the groupconsisting of Ti, V, Cr, Mn, Fe, Co and Ni, and wherein in use x rangesfrom 0 to 1. M can include Fe. The material can exhibit a solid solutionover a composition range of 0<x<0.3, or the material exhibits a stablesolid solution over a composition range of x between 0 and at leastabout 0.15, or the material exhibits a stable solid solution over acomposition range of x between 0 and at least about 0.07 or between 0and at least about 0.05 at room temperature (22-25° C.). The materialcan also exhibit a stable solid solution at low lithium content; e.g.,where 0.8<x<1 or where 0.9<x<1, or where 0.95<x<1.

In one or more embodiments, the lithium-rich transition metal phosphatephase has the composition Li_(y)MPO₄ and the lithium-poor transitionmetal phosphate phase has the composition Li_(1-x)MPO₄, wherein0.02<y<0.2 and 0.3>x>0.02 at room temperature (22-25° C.). In one ormore embodiments, the material can exhibit a solid solution over acomposition range of 0<x<0.15 and 0.02<y<0.10.

In one or more embodiments, the solid solution of the lithium transitionmetal phosphate material occupies a fraction of the compositional rangeof lithium defined as y+x.

In one or more embodiments, the lithium transition metal phosphatematerial has an overall composition of Li_(1-x-z)M_(1-z)PO₄, where Mcomprises at least one first row transition metal selected from thegroup consisting of Ti, V, Cr, Mn, Fe, Co and Ni, where x is from 0 to 1and z can be positive or negative. M includes Fe, z is between about0.15 and −0.15. The material can exhibit a solid solution over acomposition range of 0<x<0.15, or the material exhibits a stable solidsolution over a composition range of x between 0 and at least about0.05, or the material exhibits a stable solid solution over acomposition range of x between 0 and at least about 0.07 at roomtemperature (22-25° C.). The material may also exhibit a solid solutionin the lithium-poor regime, e.g., where x≧0.8, or x≧0.9, or x≧0.95.

In one or more embodiments, the lithium transition metal phosphatematerial is of a form selected from the group consisting of particles,agglomerated particles, fibers and coatings.

In one or more embodiments, the form has an average smallestcross-sectional dimension of about 75 nm or less, or about 60 nm orless, or about 45 nm or less.

In one or more embodiments, the lithium transition metal phosphatematerial is in the form of dispersed or agglomerated particles and theaverage crystallite size as determined by x-ray diffraction is less thanabout 800 nm, or less than about 600 nm, or less than about 500 nm, orless than about 300 nm.

In one or more embodiments, the form contains less than 3 wt % of asubstantially non-lithium-storing conductive phase.

In one or more embodiments, the lithium transition metal phosphatematerial is crystalline or amorphous.

In one aspect of the invention, a cathode includes a lithium transitionmetal phosphate material, for example, a lithium transition metalphosphate material having an overall composition of Li_(1-x)MPO₄, whereM comprises at least one first row transition metal selected from thegroup consisting of Ti, V, Cr, Mn, Fe, Co and Ni, and wherein in use xranges from 0 to 1. The material can exhibit a solid solution over acomposition range of 0<x<0.3 or over a range of 0<x<0.15. Anelectrochemical cell containing the electrode is also provided.

In another aspect of the invention, a nanoscale crystalline lithiumtransition metal phosphate is provided that becomes disordered upondelithiation or lithiation having a specific surface area of at leastabout 25 m²/g. In certain embodiments, a lithium deficient lithiumtransition metal phosphate is formed.

In another aspect of the invention, a lithium-deficient solid solutionlithium transition metal phosphate is provided that is formed upondelithiation at a temperature below 150° C. having a specific surfacearea of at least about 25 m²/g.

In one or more embodiments, the lithium transition metal phosphate is anordered olivine structure, and the deficiency occurs on the lithium orM1 sites of the ordered olivine, or the disorder occurs on the lithiumor M1 sites of the ordered olivine.

In another aspect of the invention, a lithium transition metal phosphateis provided that transforms upon first charge to disordered olivinehaving a lithium deficient solid solution and retains such solidsolution at temperatures below 150° C., or at temperatures below 100°C., or at temperatures below 50° C.

Still another aspect provides a high power storage battery. The batterycontains a cathode, an anode, an electrolyte in contact with andseparating the anode and cathode, a cathode current collector inelectronic communication with the cathode, and an anode currentcollector in electronic communication with the anode. The storagebattery exhibits specific power of at least about 500 W/kg (1000 W/L) atspecific energy of at least about 100 Wh/kg (205 Wh/L), and in somecases exhibits specific power of at least about 1300 W/kg (2500 W/L) atspecific energy of at least about 90 Wh/kg (180 Wh/L). In certainembodiments, the battery cathode includes a nanoscale alkalinetransition metal phosphate having a specific surface area of at leastabout 25 m²/g. In some embodiments, the cathode includes particles,fibers or coatings of a nanoscale alkaline transition metal phosphatehaving an average smallest cross-sectional dimension of about 75 nm orless. In specific embodiments, the cathode includes a composition offormula Li_(1-x)MPO₄, where M is one or more transition metals. Thecomposition has a specific surface area of at least about 25 m²/g, andexhibits a stable solid solution over a composition range of x between 0and at least about 0.03, and in some embodiments up to about 0.15. Inparticular embodiments, the cathode includes particles, fibers orcoatings of a composition of formula Li_(1-x)MPO₄, where M is one ormore transition metals. The particles, fibers or coatings have anaverage smallest cross-sectional dimension of about 75 nm or less, andthe composition exhibits a stable solid solution at room temperature(22-25° C.) over a composition range of x between 0 and at least about0.03, and in some embodiments up to 0.15.

BRIEF DESCRIPTION OF THE DRAWINGS

The following figures are presented for the purpose of illustrationonly, and are not intended to be limiting.

FIG. 1 is a transmission electron microscope image of a nanoscalelithium iron phosphate ion storage material illustrating nanoscaledimensions.

FIGS. 2A-2B show bright-field and dark-field scanning transmissionelectron microscope images, respectively, of an aggregated nanoscalelithium iron phosphate material; and FIGS. 2C-F show Fe, P, O and Celemental maps taken on the sample in FIG. 2B.

FIG. 3A is a composition-temperature phase diagram for a conventionalLi_(1-x)FePO₄ ion storage material according to certain embodiments; andFIG. 3B is a voltage vs. composition graph for a conventional orcoarsened Li_(1-x)FePO₄ material.

FIG. 4A is a composition-temperature phase diagram for a nanoscaleLi_(1-x)FePO₄ ion storage material according to certain embodiments ofthe invention demonstrating an extended region in which solid solutionis formed; and FIG. 4B is a voltage vs. composition graph for aconventional or coarsened Li_(1-x)FePO₄ material; the nanocrystallineform behaves thermodynamically and electrochemically as a distinctmaterial from the conventional or coarsened crystalline state.

FIG. 5 is a plot of discharge capacity at various C-rates for thenanoscale lithium iron phosphate of Example 2; the plot includes theinitial first charge capacity and illustrates that the first dischargecapacity is more than 10% higher than the first charge capacity.

FIG. 6 is a plot of discharge capacity at various C-rates for aconventional coarse grained lithium iron phosphate; the materialexhibits conventional first charge and discharge behavior and the plotshows a decrease in first discharge capacity compared to first chargecapacity.

FIG. 7 is a plot illustrating the equilibrium or near-equilibriumelectrical potential of a nanoscale Li_(1-x)FePO₄ ion storage materialat a nearly fully lithiated composition, according to certainembodiments, relative to a standard or reference electrode in anelectrochemical cell that allows electrochemical equilibration; anextended range of solid solution at room temperature in the nanoscalematerial is shown by a range of charge capacity, corresponding toregions of composition x, over which the open-circuit-voltage (OCV)varies continuously with composition, rather than being at a constantOCV.

FIG. 8 is a plot illustrating the equilibrium or near-equilibriumelectrical potential of a nanoscale Li_(y)FePO₄ ion storage material ata nearly fully delithiated composition, according to certainembodiments, relative to a standard or reference electrode in anelectrochemical cell that allows electrochemical equilibration; anextended range of solid solution at room temperature in the nanoscalematerial is shown by a range of charge capacity, corresponding toregions of composition x, over which the open-circuit-voltage (OCV)varies continuously with composition, rather than being at a constantOCV.

FIG. 9 shows the voltage and current traces upon charging in a PITTmeasurement of a conventional carbon-coated lithium iron phosphatesample.

FIG. 10, shows the capacity the cell of FIG. 9 at each voltage stepduring the PITT charging experiment; virtually no capacity is recordedas the voltage is raised until a large capacity is observed at theplateau voltage.

FIG. 11 illustrates a PITT discharging experiment for the cell of FIG. 9in which the first voltage step was from a charge voltage of 3.8V to avoltage that is 5 mV above the open-circuit voltage of the cell,measured at a 50% state-of-charge; virtually no discharging of the cellis seen until the PITT voltage is about 20 mV below the OCV.

FIG. 12 shows a charging PITT experiment on a nanoscale Li_(0.95)FePO₄material, in which substantial current flow, indicating charging, isseen well before the two-phase plateau voltage is reached.

FIG. 13 shows the capacity measured for the cell of FIG. 12 at eachvoltage step during the PITT charging experiment.

FIG. 14 shows a PITT discharging experiment for the cell of FIG. 12, inwhich the first voltage step was from a charge voltage of 3.8V to avoltage that is 5 mV above the open-circuit voltage of the cell,measured at a 50% state-of-charge; a substantial capacity of about 8mAH/g is measured when the PIT voltage is still 5 mV above the OCV.

FIG. 15 shows a powder X-ray diffraction pattern obtained from aconventional carbon-coated lithium iron phosphate material at 50% SOC.

FIG. 16 shows the powder X-ray diffraction pattern obtained from ananoscale LiFePO₄ sample according to the invention, measured at 67%SOC.

FIG. 17 is a schematic illustration of the spatial distribution ofspace-charge defects in a nanoscale lithium storage material accordingto certain embodiments.

FIG. 18 shows the specific capacity of the nanoscale lithium ironphosphate of Example 1 as measured from a Swagelok cell.

FIG. 19 shows test results from three lithium half-cells constructedusing Swagelok hardware as in Example 3.

DETAILED DESCRIPTION

Nanoscale ion storage materials and devices, such as storage batteries,that use these materials are provided. It has been unexpectedlydiscovered that ion storage materials having sufficiently small sizescale and correspondingly high surface to volume ratio or specificsurface area provide fundamentally different physical propertiescompared to their conventional coarse-grained counterparts. Inparticular, despite having gross structural similarities such as crystalstructure type and basic atomic arrangements, upon preparation or duringuse the nanoscale materials are compositionally and structurallydistinct from, and provide different and improved electrochemicalutility and performance compared to, the coarse-grained materials. Thedifference in relevant physical properties arises because the nanoscalematerials are sufficiently small in at least one dimension (forinstance, the diameter of an equi-axed particle, the diameter of ananorod, or the thickness of a thin film), or in 2- or 3-dimensions,that they have different defect chemical, thermodynamic, and mechanicalproperties. Nanoscale ion storage materials according to one or moreembodiments, as described herein, exhibit outstanding electrochemicalperformance for use in primary or secondary storage batteries.

In particular, the nanoscale materials provide a very high ratecapability, while providing a large fraction of the intrinsic chargecapacity and energy density of the material. The different propertiescan be exhibited, for example, in an as-prepared state, upon beingthermally equilibrated or partially thermally equilibrated (for instanceby heating), or upon equilibrating with a gas phase or condensed phasemedium, or upon being assembled and used as a bipolar electrochemicaldevice, including undergoing repeated charge-discharge cycles.

Nanoscale ion storage materials can be crystalline (i.e.,nanocrystalline) or amorphous. The unique properties discussed hereinare believed to arise from the stresses created by free or internalsurfaces or the behavior of the solid in the vicinity of a surface, andtherefore the relevant nanoscale dimension is the separation betweenfree or internal surfaces in the material. For example, for a particlethat is a single crystallite or that is amorphous, the free surfacesdefine the cross-sectional dimensions that determine the nanoscaleeffects. For a particle composed of multiple crystallites, the freesurfaces may again define the relevant cross-sectional dimensions, andif these are below the suitable size as described below, the materialwill exhibit nanoscale properties. The overall particle or aggregatesize may exceed these cross-sectional dimensions, yet a crystallitewithin the aggregate may nonetheless have cross-sectional dimensionsdefined by the separation between an internal surface (e.g., a grainboundary) and an external surface of the aggregate that are sufficientlysmall to provide nanoscale properties. Such materials will be suitablefor use in an electrochemical device wherein the crystallite hasnanoscale properties and at least a portion of the crystallite has anexternal surface that is accessible to an electrolyte phase when thenanoscale material is used in the device.

The thermodynamically, mechanically, and electrochemically distinctproperties described herein reflect a fundamental difference in natureof the nanoscale materials compared to larger scale materials, asopposed to simple or “trivial” size-scaling effects that may have beenrecognized previously in the art of battery materials. For example, therate-capability of electrode materials can be limited at least in partby solid-state diffusion of ions in the storage compound. Under suchcircumstances, an increased rate capability is expected from the use ofsmaller particles, or thinner films (in the case of thin filmbatteries), because diffusion times are shorter and charge/dischargerates correspondingly faster for a given transport coefficient ordiffusion coefficient. This simple effect of particle size is well-knownin the battery field (see, e.g., U.S. Pat. No. 5,910,382, directed toLiFePO₄ as an electrode-active material; and Zhang et al. Solid StateIonics 171:25-31 (2004), relating to LiMn₂O₄), but in no way suggeststhat other physical properties of a reduced-scale material wouldfundamentally change at a certain size scale.

As another example, transport in electrochemical systems can be limitedby surface reaction rates. A material having finer particle size andcorresponding higher surface area will naturally have higher areaavailable for surface reaction. This simple relationship again does notsuggest a fundamental change in physical properties occurring at aparticular size scale. However, the surface or interfacial chemistry ofsmall scale materials can change due to their size, potentially causinga fundamental improvement in surface reaction rate that benefits ratecapability apart from simple changes in available surface area. (See,e.g., Chiang, “Introduction and Overview: Physical Properties ofNanostructured Materials,” J. Electroceramics, 1:205 (1997), for adiscussion of unexpected differences between nanoscale materials andtheir coarse counterparts, as opposed expected differences based onwell-known size-scaling laws.)

As described in more detail below, we have discovered unique behaviorand phase composition at the nanoscale for ion storage materials basedon alkali transition metal phosphates. Examples include nanoscaleordered or partially disordered structures of the olivine (A_(x)MPO₄),NASICON (A_(x)(M′,M″)₂(PO₄)₃), VOPO₄, LiVPO₄F, LiFe(P₂O₇) or Fe₄(P₂O₇)₃structure types, wherein A is an alkali ion, and M, M′ and M″ aremetals. Many such compounds have relatively low electronic conductivityand alkali ion conductivity when conventionally prepared, such that forelectrochemical applications they benefit from unique properties arisingfrom being in the nanoscale state.

In one or more embodiments, the nanoscale ion storage material has theformula LiMPO₄, where M is one or more transition metals. In certainembodiments, the nanoscale material is an ordered olivine(Li_(1-x)MXO₄), where M is one or more of V, Cr, Mn, Fe, Co and Ni, andx can range from zero to one, during lithium insertion and deinsertionreactions. In the as-prepared state, x is typically about one. Inparticular embodiments, the special properties of nanoscale ion storagematerials may be augmented by doping with foreign ions, such as metalsor anions. Such materials are expected to exhibit similar behavior tothat demonstrated herein for Li_(1-x)FePO₄ at the nanoscale, based onthe scientific principles underlying such behavior. However, doping isnot required for a material to exhibit special properties at thenanoscale.

In other embodiments, there is some substitution of Li onto the M-site.In one embodiment, there is about 5 or 10% substitution of Li onto theFe site. The lithium transition metal phosphate material has an overallcomposition of Li_(1-x-z)M_(1-z)PO₄, where M comprises at least onefirst row transition metal selected from the group consisting of Ti, V,Cr, Mn, Fe, Co and Ni, where x is from 0 to 1 and z can be positive ornegative. M includes Fe, z is between about 0.15 and −0.15. The materialcan exhibit a solid solution over a composition range of 0<x<0.15.

FIG. 1 is a transmission electron microscope image of a nanoscalelithium iron phosphate ion storage material exhibiting particledimensions on these scales. FIGS. 2A and 2B show bright-field anddark-field scanning transmission electron microscope images,respectively, of an aggregated nanoscale lithium iron phosphatematerial. FIGS. 2C-2F show Fe, P, O and C elemental maps taken on thesample in FIG. 2A, showing that the distribution of these elements isuniform, i.e. that there are not distinguishable phases or particlesrich in one or another of these main constituents.

These nanocrystalline form compositions will possess measurably distinctproperties as described herein compared to their larger scalecounterparts. For example, the nanoscale materials retain a greaterextent of solid solution nonstoichiometry, namely, retain a higherdefect content than the coarse-grained material. Such properties aremeasurable by electrochemical and crystallographic methods well-known tothose skilled in the art. When used in electrodes for practicalapplications, such as a storage battery or other electrochemical device,the nanoscale ion storage materials provide higher charge storage athigher rates of charge or discharge than comparable materials that arenot nanoscale.

The nanoscale dimensions that realize the benefits as described hereincan be characterized by several methods. Based on results as describedin the Examples below, the size-dependent nonstoichiometry and relatedbeneficial properties of nanoscale LiFePO₄ and other ion storagecompounds increase as the particle size decreases. These properties aresignificant, measurable, and beneficial at particle sizes below thatcorresponding to a BET specific surface area of about 20 m²/g. In someinstances, materials having a BET specific surface area of at leastabout 25 m²/g, for example, at least about 30 m²/g, at least about 35m²/g, at least about 40 m²/g, at least about 45 m²/g, or at least about50 m²/g are employed. As used herein, “the BET method” refers to themethod of Brunauer, Emmett and Teller, well-known to those skilled inthe art of powder characterization, in which a gas phase molecule (suchas N₂) is condensed onto the surfaces of a material at a temperature(such as 77 K) where the coverage of condensed gas per unit area iswell-known, and the total amount of condensed gas on the sample is thenmeasured upon being liberated by heating.

For a given value of the BET specific surface area, and knowing thespecific gravity of the material, it is also possible to calculate acorresponding “equivalent spherical particle size.” This is the particlediameter that would result in the measured surface area if the materialwere in the form of identically-sized spherical particles, and is a goodapproximation of the number-averaged or mean particle size if theparticle shape is equi-axed. The particle morphology of thenanomaterials described in certain of the Examples below is nearlyspherical, and the equivalent spherical particle size calculated fromthe BET specific surface area is very close to the average particlediameter directly observed by electron microscopy. Furthermore, the sizeof crystallites or primary particles, when the materials of theinvention are crystalline, can be determined by X-ray line-broadeningmethods well-known to those skilled in the art. Thus, in certainembodiments, the nanomaterials described herein have an average (i.e.,mean) diameter of about 100 nm or less. In some instances, the averagediameter is about 75 nm or less, for example, about 70 nm or less, about60 nm or less, about 50 nm or less, about 45 nm or less, about 40 nm orless, or about 35 nm or less.

The unique properties of a nanomaterial may depend on the smallestcross-sectional dimension. Cross-sectional dimension is here understoodto be that family of straight lines that can be drawn through the centerof mass of an isolated or separable object. By assuming sphericalmorphology, the equivalent spherical particle size gives the largestaverage cross-sectional dimension of a particulate material. On theother hand, a very thin but continuous film, or a very thin butcontinuous fiber, can exhibit nanoscale effects, even though thedimensions are far larger than nanoscale in the plane of the film oralong the axis of the fiber. However, if the smallest cross-sectionaldimension, namely the thickness of the film or the diameter of thefiber, is sufficiently small, nanoscale properties may be obtained.Thus, in certain embodiments, for anisometric particles, such asnanorods, nanoplatelets, nanofibers or continuous thin films, thespecific surface area and the equivalent spherical particle size may notadequately define the characteristic dimension below which thenanomaterial will exhibit special properties. That is, for highlyanisometric particle shapes, in some instances the BET surface area canbe larger than the above-mentioned values, yet the material still willexhibit a smallest characteristic dimension sufficiently small toexhibit nanoscale properties as described herein.

If particle morphology is well-known and uniform amongst particles in asample (for instance, if the average size and aspect ratio of nanorodsor nanoplatelets is known, or even if the distribution of suchparameters is known), a specific surface area above which nanoscalebehavior will be observed can be computed for a given particle shape.However, for simplicity, in at least some such embodiments, nanoscalebehavior will be observed if the primary particles of the powder exhibita smallest cross-sectional dimension that is, on a number-averaged basisto provide a mean value, about 100 nm or less. In some instances, thesmallest cross-sectional dimension about 75 nm or less, for example,about 70 nm or less, about 60 nm or less, about 50 nm or less, about 45nm or less, about 40 nm or less, or about 35 nm or less. Thesedimensions can be measured using various methods, including directmeasurement with an electron microscope of the transmission orsecondary-electron type, or with atomic force microscopy. Herein, aprimary particle dimension is considered to be the characteristicspatial dimension that a BET surface area measurement would interrogateby adsorbing gas onto exposed surfaces of the material. In the instanceof a substantially fully-dense polycrystalline aggregate, it is thedimension of that aggregate. In the case of well-dispersed individualcrystallites, it is the crystallite dimension. In the case of particlesjoined into a sintered network or a porous assembly of the particles, itis the cross-sectional thickness of the branches of the network, or themean separation between pores that are open to the exterior of theassembly. In the case of an aggregated powder, the agglomerate may havean average crystallite size of less than about 800 nm, or less thanabout 600 nm, or less than about 500 nm, or less than about 300 nm. Insome embodiments, the nanoscale material is a thin film or coating,including a coating on a particle of any size, in which the film orcoating has an average thickness of about 100 nm or less, in some casesabout 75 nm or less, for example, about 70 nm or less, about 60 nm orless, about 50 nm or less, about 45 nm or less, about 40 nm or less, orabout 35 nm or less. The thickness of the film or coating can bemeasured by various methods including transmission electron microscopyor other microscopy methods that can view the film or coating incross-section.

In certain embodiments, the nanoscale ion storage materials describedherein are prepared from conventional materials by size-reductionprocesses (e.g., milling) to reduce the particle dimensions into thedesired range. However, this can be a highly energy-intensive process.Thus, as illustrated in the Examples below, the materials also can besynthesized in the nanoscale state, by methods including, but notlimited to, solid-state reactions between metal salts, wet-chemicalmethods, such as co-precipitation, spray-pyrolysis, mechanochemicalreactions, or combinations thereof. Nanoscale materials with the desiredparticle sizes and specific surface areas are obtained by usinghomogeneous reactants, minimizing the reaction or crystallizationtemperature (in order to avoid particle coarsening), and avoidingformation of liquid phases in which the product is highly soluble (whichalso tends to lead to particle coarsening). Specific processingconditions can typically be established for a given process withoutundue experimentation by those skilled in the art.

In some embodiments, nanoscale ion storage materials are prepared bynon-equilibrium, moderate temperature techniques, such as wet-chemicalor low temperature solid-state reactions or thermochemical methods. Thematerials thus prepared can acquire properties such as increasednonstoichiometry and disorder and increased solubility for dopantsbecause they are synthesized in a metastable state or because kineticpathways to the final product differ from those in conventional hightemperature processes. Such disorder in the nanoscale form can also bepreserved substantially under electrochemical use conditions and providebenefits as described herein.

Until the present experimental results were obtained, it was not knownif nanoscale ion storage materials would exhibit fundamentally differentphysical properties compared to their coarse-grained counterparts, norwas it known what measurable physical properties would differ, nor thesize scale that would realize these differences. Useful and advantageouscharacteristics of nanoscale ion storage materials according to certainembodiments include, but are not limited to, the following.

The materials can exhibit increased electronic conductivity, forexample, due to the co-existence in solid solution of higherconcentrations of mixed-valence transition metal ions, or changes in theelectronic structure related to a closer separation between atomicorbitals providing higher electronic carrier mobility, or both.Typically, the improved electronic conductivity will have a valuegreater than about 10⁻⁸ S/cm.

The materials can have improved electromechanical stability, such asimproved resistance to fracture, due to suppressed or delayed phasetransformations during use as a storage electrode. This can allow higherenergy, higher rate capability, and longer life of the materials andelectrochemical cells using the materials. When electrochemical cyclingcauses phase transformations, the materials also may exhibit smallermolar volume differences between phases, which contributes to morefacile transformation between the phases upon insertion and deinsertionof lithium.

In compounds where ion diffusion has reduced dimensionality, forexample, being one-dimensional (along channels) or two-dimensional(along planes) in the crystal structure, the nanoscale material canexhibit increased rate of intercalation, due to the existence ofmultiple paths out of the particle when there may be blocking immobileions in the diffusion paths. The diffusion coefficient should be amaterials property, not size dependent unless something else changessuch as structure or disorder. This phenomenon is illustrated asfollows. A particle that is 100 unit cells wide in spatial dimension,assuming each unit cell contains one formula unit of the compound, canhave 1% disorder and have only, on average, one disordered atom blockinga given diffusion channel. This will have little impact on diffusion ofions into and out of the particle, since the diffusion channel can beaccessed from both ends. In contrast, for a much larger particle havingthe same degree of disorder, the blocking ions will prevent access tothe majority of the channel. The specific value of the chemicaldiffusion coefficient of the transported ion (e.g., Li in a lithiumbattery) can be improved by the additional disorder of a nanoscalematerial, typically to a value greater than about 10⁻¹⁶ cm²/sec.

These observed properties provide an ion storage material with increasedcharge storage at higher charge and discharge rates.

Nanoscale ion storage materials as described herein differ from theirlarger scale counterparts in the composition range in which they canstably exist. In at least some embodiments, the nanoscale compound canexist in a state of extended solid solution compared to thecoarse-grained compound at the same temperature. The existence ofsolid-solution nonstoichiometry is important for improving ion andelectron transport, as has been demonstrated in numerousion-intercalation compounds

One aspect of the invention provides a nanocrystalline compositionexhibiting a much wider range of solid solution or defect content at agiven temperature than a bulk crystal or coarse powder of nominallysimilar composition and crystalline phase before phase-separating intotwo or more phases. These features are described in particular detailfor Li_(1-x)FePO₄, however, it will be apparent to those of skill in theart that application of these principals to other ion storage materialswill provide similar results.

As a non-limiting example, the conventional compound Li_(1-x)FePO₄ isknown to exhibit negligible solid solution nonstoichiometry x at roomtemperature, x being about 0.002 according to some published literature(Delacourt et al., “Two-phase vs. one-phase Li⁺ extraction/insertionmechanisms in olivine-type materials,” Abstract 200, 207^(th) Meeting ofThe Electrochemical Society, Quebec City, Calif., May 15-20, 2005;Delacourt et al., “The existence of a temperature-driven solid solutionin Li_(x)FePO₄ for 0≦x≧1,” Nature Materials, 4:254-260 (2005)), about0.0475 in another publication (V. Srinivasan and J. Newman, Journal ofthe Electrochemical Society, 151:A1517-A1529 (2004), and about 0.038 inanother publication (A. Yamada, H. Koizumi, N. Sonoyama and R. Kanno,Electrochemical and Solid State Letters, 8:A409-A413 (2005). Theconcentration of lithium that is tolerated in the delithiated compoundLi_(y)FePO₄, with which Li_(1-x)FePO₄ coexists, is even less. Thesefeatures are illustrated in the composition-temperature phase diagramfor LiFePO₄—FePO₄, shown in FIG. 3A. The phase composition for an ironphosphate with varying levels of lithium will vary with temperature, anda solid solution exists over wider ranges of lithium concentration atelevated temperatures, e.g., above 150° C. Elevated temperatures are notpractical for most ion storage applications and practical applicationsare constrained to be only slightly elevated above room temperature,e.g., less than about 100° C. Unless otherwise stated, we refer tocompositions at a temperature below about 100° C. and typically at roomtemperature (22-25° C.).

The phase diagram in FIG. 3A shows that at this temperature range, thesolid solution ranges are extremely limited. An illustrative voltage vs.composition plot at room temperature for the ion storage material isshown in FIG. 3B and demonstrates that the voltage curve is flat overmost of the compositional range, indicating the presence of a two phasesystem over almost the entire lithium composition range. InLi_(1-x)FePO₄ of a conventional coarse-grained form, the absence ofsolid solution nonstoichiometry is manifested by decomposition oflithium deficient compositions into two highly stoichiometric compoundshaving a chemical composition approach that of the end groupcompositions, LiFePO₄ and FePO₄. Both of these compounds have lowelectronic conductivity, due at least in part to the existence of nearlya single iron valence state, Fe²⁺ and Fe³⁺ respectively, in theindividual crystallites. In nearly stoichiometric LiFePO₄, the lithiumdiffusion coefficient is likely also very low, due to the absence oflattice vacancies to facilitate Li transport.

In contrast, nanocrystalline Li_(1-x)FePO₄ and Li_(y)FePO₄ having aspecific surface area measured by the BET method of greater than about20 m²/g, and in some instances greater than about 30 m²/g, has beenfound to exhibit x (and y) that are several fold larger than in theconventional compound. Indeed at room temperature, Li_(1-x)FePO₄ canexhibit x as large as 0.05, 0.07, 0.10, 0.15, 0.3 or even greater and ycan be as large as 0.05 or 0.1 or 0.2. As illustrated in FIG. 4A, thedashed lines shows the existence of a significant solid solution attemperatures of less than about 50° C. for Li_(1-x)FePO₄ andLi_(y)FePO₄. An illustrative voltage vs. composition plot at roomtemperature for the ion storage material is shown in FIG. 4B. The curvehas a demonstrably smaller flat region, indicating that thecompositional range of a two-phase system is limited. The slopingregions that flank the flat region indicate the existence of a solidsolution. The solid solution end-limits for the two end member phasescoexisting in the nanoscale material are larger than for theconventional material. For instance, in LiFePO₄ this means having alarge lithium deficiency x in the lithium-rich Li_(1-x)FePO₄ endmember,and a large lithium excess y in the lithium-deficient endmemberLi_(y)FePO₄, the ideal limiting compositions of these two co-existingphases being LiFePO₄ and FePO₄ respectively. Thus, duringelectrochemical cycling, the co-existing phases include a large extentof nonstoichiometry. The higher degree of nonstoichimetry indicates agreater population of both Fe²⁺ and Fe³⁺ at every point within thetwo-phase region, which provides higher electronic conductivity for thematerial. In addition, the sloping voltage curve of the nanophosphatepermits the functional advantage of allowing state-of-charge monitoringthat is not possible or is more difficult and expensive to conduct withmaterials exhibiting a flat two-phase discharge voltage profiles.

Improved electron and ion transport rates are well-known to improve therate capability of ion storage materials used in battery technology. Incertain lithium transition metal phosphate compounds described herein,both electron and ion transport rates are slow compared to that in somepreviously used materials (such as LiCoO₂ or LiMn₂O₄), and those skilledin the art have sought methods by which such transport can be improved.Nanoscale lithium transition metal phosphate compounds exhibit retentionof solid-solutions at various states of charge (lithium concentration),and the resulting materials exhibit high rate capability and high energythat have not previously been attainable in these materials.

The nonstoichiometry of non-alkali elements in the subject materialsalso can vary in the nanocrystalline form. The fundamentally differentphase behavior applies to each of the components of the compositionalsystem, although likely to different degrees. Other aspects of atomiclevel disorder likely also are affected at nanoscale dimensions. Forexample, in Li_(1-x)FePO₄, the site occupancy of the M1 and M2 sites ofthe ordered olivine structure, occupied solely by Li and Fe in the idealcrystal, can vary in the nanoscale material. There can be disorder ormixing of the Li and Fe cations between the two sites, and vacancydefects can appear on one or both sites. Also, solute cations (dopants)can be more soluble in the nanocrystalline material, or can occupydifferent sites than they do in the conventional material. In thenanocrystalline state, nonstoichiometry on the oxygen sublattice of thecrystal structure also can occur. The solubility of foreign anions, suchas sulfur or halogens, can increase as well. In certain embodiments,nanoscale ion storage materials as described herein exhibit one or moreof these variations in defect or solid solution behavior. However, asshown by experimental results presented herein, the presence of foreignmetals or anions is not necessary to create or define the specialproperties of the nanocrystalline state.

Differences in physical properties exhibited by the nanoscale materialsaccording to one or more embodiments of the invention compared to theirconventional coarse-grained counterparts are readily measurable bystandard thermal and electrochemical techniques, such as calorimetry,cyclic voltammetry, galvanostatic intermittent titration (GITT), orpotentiostatic intermittent titration (PITT). The improved performanceof the nanoscale materials in ion storage applications is also readilymeasurable, for example, by formulating the nanoscale material into anelectrode coating, constructing a nonaqueous electrochemical cell, andperforming charge-discharge tests at various current rates.

The state of extended solid solution in the nanoscale material can beconfirmed using electrochemical methods. For example, a compound ofnanocrystalline Li_(1-x)FePO₄ can be tested in a nonaqueouselectrochemical cell. The nanocrystalline Li_(1-x)FePO₄ serves as thepositive electrode against a source of lithium having a total lithiumcontent much greater than the lithium storage capacity of thenanocrystalline electrode, such as lithium foil. This electrochemicalcell construction is often referred to as a lithium half-cell by thoseskilled in the art of lithium-ion batteries. In such a cell, thenanoscale ion storage material is formulated into an electrode,typically using a conductive additive, such as carbon, and a polymericbinder. The nanoscale ion storage material electrode is separated fromthe lithium metal counterelectrode, typically by a microporous polymerseparator. The cell is then infused with a nonaqueous lithium-conductingliquid electrolyte. The charge and discharge rates of the electrode aresufficiently fast that the electrochemical behavior of the nanoscalematerial can be tested.

The existence of solid solution lithium deficiency is detectable as theappearance of a smaller total Li content that can be extracted from thenanocrystalline electrode upon first charging the cell, than can bere-inserted into the electrode upon discharging the cell. Thisdifference in first-charge capacity compared to first-discharge orsubsequent discharge capacity reveals the existence of lithiumdeficiency in the nanocrystalline material in its as-prepared state, andupon being assembled into a working cell. The extractable lithium isless than the amount of lithium that the same electrode can take up atsaturation. FIG. 5 illustrates this behavior for a nanoscale lithiumiron phosphate having the composition Li_(0.99)FePO₄ (Example 2). Theinitial data records the first charge capacity; subsequent data recordsdischarge capacity at different c-rates. Note that first discharge atC/5 rate is more than 11% greater than initial capacity. Note alsothat >90% discharge capacity is maintained up to 10C, which represents aremarkably high capacity at high discharge rates. The tests areconducted at a sufficiently slow rate upon both charge and discharge,and over a similar voltage range, that the observed results reflect thecapabilities of the storage material itself, rather than polarization orkinetic limitations due to the cell construction. Methods to ensure thatsuch is the case are well-known to those skilled in the art.

This behavior observed for nanocrystalline lithium iron phosphate isstrikingly different from that of a conventional or coarse LiFePO₄, andindeed that of most insertion electrode materials. Such materialstypically exhibit a first-charge capacity that is greater than thefirst- and subsequent discharge capacities using a similar cellconfiguration. Results from one comparative example are shown in FIG. 6.Comparison of this conventional material to the nanoscale material inFIG. 5 highlights some striking differences. First, discharge capacityat C/5 decreases by more than 10% from first charge capacity and thedischarge capacity decreases steadily with increasing discharge rate.

The advantages imparted by a nanoscale material according to one or moreembodiments of the present invention are counter-intuitive because ahigh, initial charge capacity is typically associated with a greaterextractable lithium content. While it is generally desirable for alithiated electrode material to have a higher initial extractablelithium content, in the present instance the ability of the nanoscalematerial to sustain a lithium-deficient solid solution confers variousadvantages as described herein, which may overcome the disadvantage ofhaving slightly less lithium capacity.

Moreover, as discussed later, the nanoscale materials of the inventioncan sustain a nonstoichiometry x and y in the coexisting phase that maybe as large or larger than the nonstoichiometry present in the as-madematerial. Thus preparation in an initially nonstoichiometric state isnot required of the materials of the invention, nor necessary in orderto obtain the benefits described herein.

The ability of nanocrystalline materials as described herein to exist ina more highly nonstoichiometric or defective state than their coarsecounterparts also can be demonstrated by measuring the equilibrium ornear-equilibrium electrical potential of a nanoscale material relativeto a standard or reference electrode in an electrochemical cell thatallows electrochemical equilibration. It is well-known to those skilledin the art that the equilibrium electrical potential of such a cell,having one electrode whose potential is suitably well-referenced, can beused to determine the chemical potential of an electroactive species inthe other, test electrode.

FIG. 7 shows the cell voltage vs. specific capacity of the positiveelectrode active material for cells in which a lithium metalcounterelectrode has been used, serving as a suitable reference. Twonanoscale lithium iron phosphate materials of overall compositionsLiFePO₄ and Li_(0.95)FePO₄ are compared against a conventional,commercially available carbon-coated lithium iron phosphate. All threecells are tested at a slow C/50 rate permitting the near-equilibriumcell voltage to be observed. The nanoscale materials are further knownfrom separate tests to exhibit much faster relaxation to theirequilibrium potentials than does the conventional sample. It is seenthat the nanoscale materials exhibit a substantial charge capacity overwhich the voltage varies continuously, before reaching a relativelyconstant voltage plateau. In contrast, the cell voltage for theconventional material exhibits no such regime, instead reaching itsvoltage plateau nearly immediately after a small voltage overshoot.

FIG. 8 shows the C/50 discharge curves for the same three samples. Hereit is seen that at the beginning of discharge the nanoscale materialsboth exhibit a capacity regime of continuously varying voltage,indicating the existence of a solid solution, that is essentially absentfor the conventional material, and at the end of discharge, bothnanoscale materials exhibit a wide capacity regime of continuouslyvarying voltage indicating a solid solution. These examples demonstratethe effect pictorially illustrated in FIGS. 3B and 4B for nanoscale andconventional lithium iron phosphate materials, respectively.

Other accepted electrochemical methods that can be used to show that thenanoscale materials of the invention possess regimes of extended solidsolution include GITT and PITT. In GITT, the open-circuit-voltage (OCV)measured after allowing an electrochemical cell to approach equilibriumwill exhibit a composition dependence (i.e., as a function ofstate-of-charge or charge capacity) that is measurably different betweenthe conventional and nanocrystalline forms. An extended range of solidsolution in the nanoscale material is shown by regions of composition xover which the OCV varies continuously with composition, rather thanbeing at a constant OCV. This indicates a constant chemical potentialfor lithium despite variation of x, corresponding to a multi-phaseequilibrium. Such measurements typically can be conducted to ±0.002V orbetter precision by those skilled in the art, allowing comparison ofdifferent materials to determine the value of x at which the boundarybetween a single-phase solid solution and multiple phases lies. For ananoscale material, there is a wider range of composition x over whichthe single-phase solid solution can exist. The wider range of solidsolution in the nanoscale form can be attained for any one or more ofthe individual phases exhibited by the compound, including intermediatephases forming within the limits of lithiation discussed here.

The PITT method is also useful for not only determining the cellvoltages at which electrochemical oxidation and reduction of anelectrode-active compound occur, but also for providing informationregarding the rate and mechanism of such reactions. In PITT, the cellvoltage is stepped upwards or downwards incrementally, and the currentflow is monitored as the cell spontaneously charges or discharges. FIG.9 shows the voltage and current traces upon charging in a PITTmeasurement of a conventional carbon-coated lithium iron phosphatesample. With each incremental voltage step of 10 mV, the current isobserved to flow as the cell undergoes charging. It is notable thatvirtually no capacity is recorded until a voltage plateau is reached.Also, during charging on the voltage plateau, the current flow risesslowly over a period of several hours and then decays, showing sluggishkinetics for the phase transformation occurring during charging. In FIG.10, the capacity measured for the cell at each voltage step during thePITT charging experiment is shown. It is seen that virtually no capacityis recorded as the voltage is raised until a large capacity is observedat the plateau voltage. In FIG. 11 are shown results for the same cellduring a PITT discharging experiment in which the first voltage step wasfrom a charge voltage of 3.8V to a voltage that is 5 mV above theopen-circuit voltage of the cell, measured at a 50% state-of-charge. Inthis experiment, virtually no discharging of the cell is seen until thePITT voltage is about 20 mV below the OCV.

The nanoscale materials of the invention behave in markedly differentmanner. FIG. 12 shows a charging PITT experiment on a nanoscaleLi_(0.95)FePO₄ material, in which substantial current flow, indicatingcharging, is seen well before the two-phase plateau voltage is reached.In addition, with each upward voltage step, the maximum in current isobserved immediately, rather than several hours into the current decayprocess as seen in FIG. 9. This shows that the phase transformationforming the delithiated Li_(y)FePO₄ phase is more facile in thenanoscale material. FIG. 13 shows the capacity measured for the cell ateach voltage step during the PITT charging experiment. It is seen thatthere is substantial charging occurring below the plateau voltage. Notethat because charging can only occur when the applied voltage is equalto or greater than the equilibrium voltage, this result shows that thereexist compositions with an equilibrium voltage below that of thetwo-phase plateau. That is, it demonstrates the existence of alithium-deficient solid solution Li_(1-x)FePO4. In FIG. 14 are shownresults for this same cell during a PITT discharging experiment in whichthe first voltage step was from a charge voltage of 3.8V to a voltagethat is 5 mV above the open-circuit voltage of the cell, measured at a50% state-of-charge. Here, a substantial capacity of about 8 mAh/g ismeasured when the PITT voltage is still 5 mV above the OCV. Since upondischarge, no driving force exists until the applied voltage is at orbelow the equilibrium voltage, this result demonstrates the existence ofa lithium excess solid solution Li_(y)FePO₄ at voltages above theplateau voltage.

The differences between nanoscale Li_(1-x)FePO₄/Li_(y)FePO₄ and theconventional materials can also be quantified by X-ray diffraction. Thepresence of a compositionally distinct nonstoichiometry in nanoscaleLi_(1-x)FePO₄ is demonstrated by unique lattice constants (a, b and cwithin the orthorhombic unit cell) and unique unit cell volume (given bythe product a×b×c). Conventional crystalline olivine LiFePO₄ has alarger a and b lattice parameter, and a smaller c lattice parameter,than does crystalline FePO₄. A continuous solid solution between LiFePO₄and FePO₄ would therefore show a continuous variation between thelimiting values of the lattice constants as the lithium concentrationvaries between one and zero. The lattice constants of the materialsaccording to one or more embodiments of the invention may therefore beused to determine the corresponding nonstoichiometry of the coexistingphases. This was accomplished by carrying out careful X-ray diffractionmeasurement of the subject materials at different states of lithiation(different states of charge, SOC), from which lattice parameters andother crystallographic information was obtained using Rietveldrefinement, a process for analyzing diffraction data that is well-knownto those skilled in the art of battery materials synthesis andcharacterization.

FIG. 15 shows a powder X-ray diffraction pattern obtained from aconventional carbon-coated lithium iron phosphate material (AldrichChemical) at 50% SOC. To this sample was added silicon powder to providean internal standard for the X-ray peak positions. It is seen that thepeaks for LiFePO₄ olivine are well aligned with the expected peakpositions for this phase, based on the data in reference 01-081-1173from the Joint Committee on Powder Diffraction Standards (JCPDS). Thepeaks for the olivine form of FePO₄ are also seen in FIG. 15, and aresomewhat displaced from the positions for a somewhat differentcomposition listed by JCPDS.

FIG. 16 shows the powder X-ray diffraction pattern obtained from ananoscale LiFePO₄ sample according to the invention, measured at 67%SOC. It can be seen that numerous peaks for both the “LiFePO₄” and“FePO₄” phases are displaced from their corresponding positions in FIG.15. A precise determination of the lattice constants in these materialswas made using the Rietveld refinement method, on powder X-raydiffraction spectra carefully obtained over a wide diffraction anglerange (known to those skilled in the art as the “2-theta” range) of 15degrees to 135 degrees. It was found that the nanoscale materialsaccording to one or more embodiments of the invention, when in a stateof charge such that the above mentioned two olivine phases co-exist,have distinctly different lattice parameter values from the conventionalmaterial. The lattice parameters and unit cell volumes are reported inTable 1, in which the nanoscale lithium iron phosphate was measured at67% state-of-charge, compared with similar measurements made for aconventional LiFePO4/FePO4 reported in the literature (A. S. Anderssonand J. O. Thomas, J. Power Sources, 97-98: 498 (2001)). For example, onthe lithium rich side of the phase diagram, nanoscale Li_(1-x)FePO₄having smaller a and b lattice constants and a larger c lattice constantfor than conventional LiFePO₄ are obtained. The lithium deficient solidsolution coexists with an Li_(y)FePO₄ phase having the latticeparameters for a, b that are larger and c that is smaller than inconventional FePO₄. These measurements show that indeed x and y arelarger than their corresponding values in conventional LiFePO₄/FePO₄,notwithstanding some smaller nonstoichiometry existing in thosematerials as well. From the Rietveld refinement of the nanoscale sample,a crystallite size of about 28 nm was determined, which is close to thecalculated equivalent spherical particle size of 36.1 nm and shows thatthe high surface area of the sample is due to nanoscale crystallites ofthe lithium iron phosphate and not due to a high surface area impurityor additive phase.

TABLE 1 Lattice constants and unit cell volume for LiFePO₄, FePO₄,Li_(1−x)PO₄ and Li_(y)FePO₄ Unit cell Material a (Å) b (Å) c (Å) volume(Å³) LiFePO₄ 10.329 6.007 4.691 291.02 Li-deficient 10.288 5.991 4.698289.56 Li_(1−x)FePO₄ FePO₄ 9.814 5.789 4.782 271.7 Li-rich Li_(y)FePO₄9.849 5.809 4.781 273.55

The existence of larger nonstoichiometry in the coexisting phases of thematerials of the invention is therefore readily measured usingdiffraction methods. The value of x and y determines the ratio of 2+ to3+ transition metal valences (in the case of iron, it is Fe²⁺/Fe³⁺) inthe materials, and larger values correspond to a higher concentration ofthe minority valence state. This has the effect of increasing theelectronic conductivity of each phase compared to the same phase in itsconventional state of lower x or y, and thereby improves electrochemicalperformance of the battery. In addition, a reduction in the latticeparameter of the Li_(1-x)FePO₄ phase (or any other compositions of thelithium-rich endmember) has the effect of bringing the multivalenttransition metal ions closer together within the structure, which alsoincreases the degree of orbital overlap thereby changing the electronicstructure of the material so as to decrease the bandgap or increasecarrier mobility, thereby increasing electronic conductivity.

The a, b lattice constants for the lithium deficient Li_(1-x)FePO₄ isless than that for LiFePO₄ and the a, b lattice constants for thelithium rich Li_(y)FePO₄ is greater than that for FePO₄. Therefore, themismatch in lattice parameters and unit cell volume is decreased in thenanoscale materials of the invention, which may have a profoundinfluence on the electrochemical performance of the material,particularly at high charge/discharge rates. This is because thefacility with which one phase is formed from the other upon charging anddischarging of the electrochemical cell is dependent on the mismatch inlattice parameters (if crystalline) and the relative volumes of the twoco-existing phases.

The lattice parameters and unit cell volume between the coexistingphases Li_(1-x)FePO₄ and Li_(y)FePO₄ and the unit cell volume arereported in Table 1. From these values, one may readily compute thedifferences in lattice parameters and in unit cell volumes, on apercentage basis, for a transformation from the Li_(1-x)FePO₄ toLi_(y)FePO₄ phase, which corresponds to charging of a cell using thelithium iron phosphate as the positive electrode, or from Li_(y)FePO₄ toLi_(1-x)FePO₄, which corresponds to discharging. The percentage changesupon charging are found to be slightly smaller than those upondischarging, and this may cause differences in the inherent rate ofcharging versus discharging within any one material. However, tofacilitate comparison of the nanoscale and conventional materials, wemay also compute the differences in the respective values of any latticeconstant or unit cell volume as a percentage of the mean value betweenthe two, as has been done in Table 2. That is, the percentage differencein the a lattice constant is the difference in a between any twomaterials divided by the arithmetic mean value of a for those twosamples. Herein, unless otherwise stated, the percentage differences arecomputed in this manner. For nanoscale Li_(y)FePO₄/Li_(1-x)FePO₄ thedifferences in lattice parameters are Δa=4.36%, Δb=3.07%, Δc=−1.75%, andthe difference in unit cell volume is ΔV=5.69%. In comparison, forconventional LiFePO₄/FePO₄, the corresponding numbers are Δa=5.11%,Δb=3.68%, Δc=−1.93%, and ΔV=6.87% for the limiting endmembers. We alsomeasured a conventional material (Aldrich Chemical) that was taken to50% state-of-charge, and in which the co-existing compositions have thesmall permitted extent of nonstoichiometry. Here the difference isΔa=4.91%, Δb=3.64%, Δc=−2.03%, and ΔV=6.52%. These unit cell and latticeparameter differences are summarized in Table 2.

Although not shown in Table 2, one may also readily compute the misfitstrain of a plane separating the two limiting compositions Li_(y)FePO₄and Li_(1-x)FePO₄. This is important because the formation of one phasefrom the other during electrochemical cycling must necessarily introducean interface between the two materials, which is a two-dimensionalfeature. Inspection of the results in Table 1 shows that the planeformed by the principal axes a and b (the ab plane or in Miller indicesthe {110} plane) have the largest difference in area between theLi_(y)FePO₄ and Li_(1-x)FePO₄. the ac plane (or {101}) has the nextlargest difference, and the bc plane (or {011} has the least difference.This indicates that the bc plane is the most preferred orientation alongwhich one phase will grow topotaxially upon the other (or vice versa).Comparing the nanoscale and conventional materials in Table 1, thesedifferences are 7.43%, 2.62% and 1.32% respectively for the nanoscalematerial, and 8.79%, 3.19%, and 1.76% respectively for the conventionalmaterial. In the Aldrich material measured at 50% SOC these differencesare 8.55%, 2.88% and 1.62% respectively. Thus, according to oneembodiment, the nanoscale materials of the invention are defined byhaving a plane formed by any of the principal axes of the crystal alongwhich the strain measured as a change in the area is less than about1.6%, or less than about 1.5%, or less than about 1.4%. According toanother embodiment, none of the planes formed by any of the principalaxes of the crystal have such a strain exceeding 8%, or 7.5%, or 6%.

TABLE 2 Lattice parameter and unit cell data Material Δ a (%) Δ b (%) Δc (%) Δ V (%) LiFePO₄ 5.11 3.68 −1.93 6.87 FePO₄ Li-deficient nano 4.363.07 −1.75 5.69 Li_(1−x)FePO₄ Li-rich nano Li_(y)FePO₄ Li_(1−x)FePO₄4.91 3.64 −2.03 6.52 (conventional) Li_(y)FePO₄ (conventional)

These differences between the nanoscale and conventional materials aresignificant, due to the fact that the elastic moduli of these inorganiccompounds are very high, e.g., on the order of 100 GPa. Small percentagedifferences in lattice parameters and unit cell volumes result in largeelastic energies if these highly stiff solids are made to accommodatethe strains without breaking apart. By engineering the nanoscalematerial of the invention to have small differences in latticeparameters and unit cell volumes between coexisting phases, not only isthe energy required to transform one phase from the other decreased, thelikelihood of mechanical fracture and defect formation during cycling,so-called “electrochemical grinding,” is minimized, leading to anexceptionally long cycle life for the materials of the invention.

It is also recognized that while there is a limiting particle size abovewhich the benefits seen in the materials of the invention are no longerrealizable, it is expected that there is virtually no practical lowerlimit to the particle sizes that may be obtained through synthesismethods known to those skilled in the art. As the particle size of thenanoscale materials of the invention decrease, the extent ofnonstoichiometry x and y under any given synthesis or test conditionsincreases, and the differences in lattice constants and unit cellvolumes between the coexisting phases decreases as well. That is,referring to FIG. 4, the boundaries of the two-phase regime move inwardin composition and down in temperature. For sufficiently fine particlesizes, a complete solid solution becomes achievable at room temperature.

The cycle life of a rechargeable battery is typically defined as thenumber of charge/discharge cycles, over a specified voltage range and ata specified current rate, over which the capacity of the batterydecreases to a certain percentage of the initial value. Conventionalcathode-active materials and rechargeable batteries using thesematerials, including LiFePO₄ olivine and its compositional derivatives,over a voltage range of about 2V to 3.8V and at a current rate of about1C, typically show a cycle life of less than 1000 cycles before thecapacity decreases to 80% of its initial value. In contrast, thematerials and devices of the invention can undergo in excess of 1000,even in excess of 2000, and in some instances in excess of 5000 cyclesbefore decreasing in capacity by this amount. At higher charge dischargerates, for example a 5C charge/discharge rate over the same voltagerange, conventional materials will typically show a cycle life of lessthan about 500 cycles before decreasing in capacity to 80% of theinitial value. In contrast, the materials and devices according to oneor more embodiments of the invention may exhibit greater than 1000 fullcharge/discharge cycles before decreasing in capacity by this amount.

Many applications of a high power battery, including but not limited tohybrid electric vehicle applications, require high rate charge/dischargepulses over a narrower range of voltage or capacity than full cycling.Under such conditions, the cycle life of the materials and devices ofthe invention can be extraordinarily long. One well-known pulse testingprotocol is the “HPPC” test defined by the United States AdvancedBattery Consortium (USABC). The materials of the invention, when used ina battery meeting the specific energy and specific power requirementsdefined by the USABC, are able to exhibit in excess of 150,000 of cyclelife before the performance of the battery falls below the defineduseful limits.

It is understood that during the dynamic process of lithiumintercalation and deintercalation, the stresses generated by thedifferences in lattice parameters can cause the unit cell parameters andcorresponding compositions x and y of the coexisting phases totemporarily deviate from their stable values. Nonetheless, upon allowingsome time for stress relaxation and local equilibration within thematerials, the above described differences between nanoscale andconventional materials are seen, thereby clearly distinguishing the twoclasses of materials from each other. The properties of the materialsincluding the lithium nonstoichiometry may not yet be at their stablestate when first preparing a material and assembling an electrochemicaldevice. In use as a reversible electrochemical device such as arechargeable battery, the behavior of the material during the very firstcycle may not be as important as the behavior during subsequent cycling.Therefore the differences in unit cell parameters and lithiumconcentrations desirably are measured after at least one fullintercalation and deintercalation cycle between the working voltagelimits of the device, and after allowing said material to rest in itsstate-of-charge for at least 12 hours. According to one or moreembodiments of the present invention the extent of solid solution ineach endmember phase may increase with electrochemical cycling, allowingthe transformation from one phase to the other to become more facilewith the use of the battery. This is manifested in, amongst otherbehavior, as a decrease in the impedance of the battery withcharge/discharge cycling.

In the materials according to one or more embodiments of the invention,the formation of one phase from the other (and vice versa) uponelectrochemical cycling is made much more facile in comparison toprevious materials by the fact that the materials are nanoscale, andbecause they have been engineered to have smaller lattice parameter andunit cell mismatch between the two co-existing phases. The advantages ofminimizing the mismatch stresses in order to permit facile phasetransitions and high rates of charge and discharge have not previouslybeen recognized in the field of battery materials.

Conventional understanding also teaches away from the use of highsurface area active materials in battery electrodes, especially on thepositive electrode side, for several reasons, such as poor safety,excessive self-discharge, rapid impedance buildup over time, orshortened cycle life at elevated temperatures, or low tap density andpacking density resulting in undesirably low energy density infabricated batteries. For example, it is well-known that the cathodeactive materials LiCoO₂ and LiNiO₂, including their solid solutions andderivatives, can create unsafe conditions in the highly charged statedue to the presence of their transition metals in the highly oxidized 4+valence state. Overcharged and/or overheated lithium ion cells usingthese cathode materials, even in conventional form, can exhibitthermal-runaway leading to fire or explosion, and it is generallyconsidered to be the case that such risks are exacerbated by the use ofhigher surface area active materials. Also, at elevated temperatures andover long operating times, lithium ion cells using these cathodematerials exhibit impedance rise due to interfacial reactions, whichlowers the power capability. Thus, the use of these materials in ananocrystalline state generally is considered unwise for both safety andlife reasons. As another example, the cathode active material LiMn₂O₄has been used in high power lithium ion batteries, but frequentlyexhibits permanent capacity loss after use or storage, related to thedissolution of manganese in the electrolyte and/or protonation of thesurface of the active material particles by residual acid in the liquidelectrolytes used in such cells. Since these effects are exacerbated inhigh surface area materials, common knowledge teaches away from the useof nanocrystalline LiMn₂O₄. These observations suggest that nanoscaleparticle sizes could be undesirable with respect to certain properties.However, using the nanoscale ion storage materials described herein,such difficulties can be overcome while retaining energy density andpower density advantages.

The surprisingly wider range of solid solution of the nanoscalematerials of the invention compared to their conventional counterpartsmay be due to stress, both the stress exerted by the highly curved freesurface combined with the surface tension of the material, and thestress induced when the two phases coexist and a region of each phaseeach exerts a stress on a region of the other phase. In addition, whilenot being bound by any particular interpretation, it is believed thatdifferences in the properties of the nanoscale ion storage materialsdescribed herein compared to their conventional larger scalecounterparts are also due to the formation of near-surface defect layersthat alter the overall defect thermodynamic state of the material. Thedifferences in physical properties and structure between the nanoscaleand conventional crystalline states can be likened to the differencebetween the crystalline and glassy forms of a single composition, whichhave such clearly different thermodynamic, structural and physicalproperties as to be considered different materials.

While not being bound by any mode or theory of operation, the followingmechanisms may provide a basis for unique properties of thenanocrystalline materials according to one or more embodiments of thepresent invention. In iono-covalent compounds having a latticediscontinuity, such as a free surface or grain boundary, due todifferences in the free energies of formation of lattice defects, thesurface can become enriched in one or more atomic species relative toothers. This gives rise to an excess surface charge, and a compensatingspace-charge layer that penetrates a short distance into the solid, thespace-charge layer being composed of charged defects. When thespace-charge defects are vacancies, the bulk of the crystal then has anoverall excess of the vacancies, namely an altered stoichiometrycompared to the ideal crystal in the absence of the surface orinterface. The space-charge phenomenon is well-established in ioniccrystals through many theoretical and experimental studies, includingpapers published by one of the inventors. (See, e.g., Y.-M. Chiang, D.P. Birnie, III, and W. D. Kingery, Physical Ceramics: Principles forCeramic Science and Engineering, Chapter 3, John Wiley & Sons (1997);Chiang et al., “Characterization of Grain Boundary Segregation in MgO,”J. Am. Ceram. Soc., 64:383-89 (1981); Ikeda et al., “Space ChargeSegregation at Grain Boundaries in Titanium Dioxide: Part I,Relationship Between Lattice Defect Chemistry and Space ChargePotential,” J. Am. Ceram. Soc., 76:2437-2446 (1993); Ikeda et al.,“Space Charge Segregation at Grain Boundaries in Titanium Dioxide: PartII, Model Experiments,” J. Am. Ceram. Soc., 76:2447-2459 (1993).) Wehave experimentally observed nonstoichiometry and extended solidsolution behavior consistent with space-charge influenced behavior ofnanocrystals in these materials. Accordingly, while not being bound byany particular theory, we address the possible origins of this behavior.

Consider a starting point of a stoichiometric LiFePO₄ olivine compoundthat is then allowed to equilibrate its free surface with itssurroundings. The surface is likely to become enriched in the ion havingthe lowest defect formation energy and/or sufficient mobility to beremoved preferentially to the surface. In LiFePO₄, this ion isenergetically and kinetically most likely to be lithium. Creation of alithium-rich surface must leave a lithium-deficient interior, in whichthe deficiency corresponds to the presence of lithium vacancies. As withother compounds exhibiting space-charge behavior, the lithium deficiencyis not likely to be distributed uniformly across the interior. Instead,the lithium vacancies may be preferentially concentrated near thesurface in a space-charge layer. The spatial extent of this layerdepends at thermal equilibrium on the defect concentration, thedielectric constant of the material, and the temperature. If the systemis not at equilibrium, the extent of the space-charge layer depends ontransport kinetics of the ions and defects as well.

The spatial distribution of defects is shown schematically in FIG. 17.The spatial extent of the space-charge layer can be of the order of oneto several nanometers. The near-surface concentration of vacancies orother defects can be many times greater than the concentration thatwould be tolerated in a bulk crystal as a solid solution, i.e., withouthaving precipitation or phase-separation. Thus, for a sufficiently smallnanoparticle, nanorod, nanofiber or thin film, the interior of theparticle has a measurably higher lithium deficiency than a conventionalparticle. Overall, the particle now behaves in a nonstoichiometricmanner, especially if the Faradaic behavior of the Li⁺ at the surfacediffers from that in the bulk. X-ray diffraction measurements andelectrochemical tests can detect these differences compared toconventional materials. Furthermore the surface lithium ions can bereacted easily by surface reactions with adjacent media such as liquidelectrolyte, or evaporated upon heating or reaction with the gas phaseas a lithium oxide or lithium carbonate species. In such instances, thenanoparticle is left more lithium-deficient than a conventional particleor crystal, yet said defects giving rise to the nonstoichiometry remainas a solid solution rather than causing the nanoparticles to form newand separate phases as in a conventional material. In the case of anear-surface enrichment of lithium vacancies, the Fe³⁺/Fe²⁺ ratio alsocan vary spatially with distance from the surface, and provide not onlygreater electronic conductivity to the particle as a whole, but agreater electronic conductivity at the surface of the particle than inthe interior.

In at least some embodiments, the nanoscale ion storage materialsdescribed herein typically contain less than about 5 weight percent, orabout 3 weight percent, of any additional phase that does notsubstantially store ions, but may provide added electrical conductivity.Such additional phases include, for example, carbon, a metal, or anintermetallic phase, such as a metal phosphide, metal carbide, metalnitride, or mixed intermetallic compound, such as metal carbide-nitrideor metal carbide-phosphide. In certain embodiments, for use as a storageelectrode, the nanoscale material typically is formulated into anelectrode by standard methods, including the addition of a few weightpercent of a polymeric binder, and less than about 10 weight percent ofa conductive additive, such as carbon. In at least some suchembodiments, the electrodes typically are coated onto one or both sidesof a metal foil, and optionally pressed to a coating thickness ofbetween about 30 micrometers and about 200 micrometers, providing acharge storage capacity of between about 0.25 mAh/cm² and about 2mAh/cm². Such electrodes can be used as the positive or negativeelectrode in a storage battery. Their performance can be evaluated, forexample, using laboratory cells of the coin-cell or so-called Swagelokcell types, in which a single layer of electrode is tested against acounterelectrode (typically lithium metal when the nanoscale material isa lithium storage material) using galvanostatic (constant current) orpotentiostatic (constant voltage) tests or some combination of the two.Under galvanostatic conditions, the current rate can be described as“C-rate,” in which the rate is C/n, and n is the number of hoursrequired for substantially complete charge or discharge of the cellbetween a selected upper and lower voltage limit.

In certain embodiments, when used as the positive electrode in a lithiumbattery, the electrodes are typically assembled into multilayerlaminated cells of wound or stacked configuration, using lithium metalor an anode-active lithium storage electrode as the negative electrode.Non-limiting examples of suitable negative electrode materials includelithium metal, carbon, an intermetallic compound, or a metal, metalloidor metal alloy that includes such lithium-active elements as Al, Ag, B,Bi, Cd, Ga, Ge, In, Pb, Sb, Si, Sn or Zn. The negative electrodematerial can be selected or designed for high rate capability. Thestorage batteries thus assembled can employ a porous electronicallyinsulating separator between the positive and negative electrodematerials, and a liquid, gel or solid polymer electrolyte. The storagebatteries can have electrode formulations and physical designs andconstructions that are developed through methods well-known to thoseskilled in the art to provide low cell impedance, so that the high ratecapability of the nanoscale ion storage material can be utilized.

The nanoscale ion storage materials described herein, when tested insuch laboratory cells or in storage batteries, will exhibit greatlyimproved capacity retention at high charge and discharge rates comparedto their coarse-grained counterparts. Typically, over a voltage range inwhich the upper voltage limit is about 120% of, and the lower voltagelimit is about 50% of, the average voltage exhibited by the cell at alow rate of C/5 or less, the discharge capacity measured at a 5C ratecompared to the capacity measured at a low rate of C/5 or less (i.e.,the capacity retention) will be about 80% or greater, in some casesabout 90% or greater, or about 95% or greater. At a 10C rate, thecapacity retention can be about 75% or greater, in some cases about 85%or greater, for example, about 90% or greater, or about 93% or greater.At a 20C rate, the capacity retention can be about 60% or greater, insome cases about 70% or greater, for example, about 80% or greater, orabout 85% or greater. At a 35C rate, the capacity retention can be about50% or greater, in some cases about 60% or greater, for example, about75% or greater, or about 80% or greater. At a 50C rate, the capacityretention can be about 30% or greater, in some cases about 40% orgreater, for example, about 50% or greater, or about 60% or greater.

In some embodiments, when used in a complete wound or stacked multilayercell having at least 5 Wh energy at a C/5 or lower discharge rate, thenanoscale materials described herein can provide cells with thefollowing levels of specific power (power density) and specific energy(energy density) for substantially complete discharge starting from afully charged state (i.e., 100% depth of discharge). The cells canexhibit, for example, specific power of at least about 500 W/kg (1000W/L) at specific energy of at least about 100 Wh/kg (205 Wh/L), specificpower of at least about 950 W/kg (2000 W/L) at specific energy of atleast about 95 Wh/kg (190 Wh/L), specific power of at least about 1300W/kg (2500 W/L) at specific energy of at least about 90 Wh/kg (180Wh/L), and specific power of at least about 1600 W/kg (3200 W/L) atspecific energy of at least about 85 Wh/kg (175 Wh/L). It is understoodthat for shallower depth of discharge, the specific power and powerdensity can be significantly higher than those given above.

The following non-limiting examples further illustrate certainembodiments.

Example 1

Lithium iron phosphate of overall composition LiFePO₄ was prepared usingthe following proportions of starting materials:

Li₂CO₃ (Alfa-Aesar, 99.999%) 0.739 g Iron (II) oxalate (Alfa-Aesar,99.999%) 3.598 g Ammonium phosphate (Aldrich, 99.998%) 2.301 g

While these basic components are known as starting materials for thesynthesis of conventional LiFePO₄, here through the use of a high purityacetone as the solvent (reagent grade, J. T. Baker), and using extendedmixing to allow the starting components to undergo a gas-evolvingmechanochemical reaction, a precursor is obtained that upon firingyields a low carbon, very high specific surface area nanoscalephosphate. The dry components were weighed and mixed with a sufficientquantity of high purity acetone to create a free-flowing suspension, andthe mixture was roller-milled in a sealed polypropylene jar usingzirconia milling media for 24 hours, obtaining a homogeneous andfinely-divided precursor suspension. The precursor was thoroughly driedand then heat treated in a tube furnace under flowing argon gas (grade5.0), first at 350° C. for 10 h and then at 600° C. for 20 h. After heattreatment, the specific surface area was measured using the BET methodand found to be 38.6 m²/g, for which the equivalent spherical particlediameter was calculated to be 43.2 nm, assuming a crystal density of 3.6g/cm³. The carbon content was analyzed by the combustion method andfound to be below 3 weight percent, such that the measured surface areacan be predominantly attributed to the nanoscale phosphate phase. Forpowders prepared by this procedure, transmission electron microscopyimaging such as in FIGS. 1 and 2 showed that the observed averageparticle diameter was close to the equivalent spherical particle sizecalculated from the BET specific surface area.

The fired powder was formulated into an electrode having the followingcomposition:

Nanoscale lithium iron phosphate powder 3.95 g Super P carbon 0.50 gKynar 2801 binder 0.55 g γ-butyrolactone (solvent) 28.5 gand mixed to create a free flowing suspension, then cast in a uniformlayer onto aluminum foil. The coating was dried in vacuum at 100-110°C., after which it was measured to have a thickness of about 100micrometers, and punched into discs of 1-2 cm diameter as appropriate tofit Swagelok or coin cells. The electrode coatings were assembled intolithium half-cells using Swagelok or coin cell hardware, using amicroporous polymer separator, lithium foil as the negative electrode(total lithium content at least ten times greater than the theoreticalstorage capacity of the positive electrode), and a conventionalnonaqueous lithium ion battery electrolyte containing LiPF₆ as thelithium salt. FIG. 18 shows the specific capacity of the nanoscalelithium iron phosphate as measured from a Swagelok cell. The ability ofthe nanoscale material to deliver high capacities at high charge ordischarge rates is remarkable. The discharge capacity retention here isused to describe the percentage of the capacity measured at a particularC-rate, over the voltage range 2.0-3.8V, compared to the capacityobserved at C/5 rate over the same voltage range, as shown in FIG. 16.At 1.8C rate, the capacity retention was 95.9%; at 4.4C rate, theretention was 92.1%; at 9C rate, the retention was 88.1%; at 18C rate,it was 82.6%; at 31C rate, it was 75.6%; and at 44C rate, it was 69.1%.Those skilled in the art of battery materials will recognize that theseare extraordinarily high capacity retention values compared toconventional ion storage materials. The capacity measured during thefirst charging cycle for this sample was about 6.6% less than thecapacity during the first discharge cycle, showing that thenonstoichiometry x of the as-produced material is about 6.6%.

Example 2

A nanoscale ion storage material having overall compositionLi_(0.99)FePO₄ was synthesized and tested following procedures asdescribed in Example 1, except that a larger batch size was made anddifferent sources of starting materials were used. The composition wasmade using the following proportions of starting materials:

Li₂CO₃ (SQM)  7.4337 g Iron (II) oxalate (Elementis) 36.2696 g Ammoniumphosphate (Heico) 22.5541 g

A larger sealed polypropylene container, and steel milling media wereused to mill the starting materials for 72 hours. Firing of the driedpowder was conducted in nitrogen of 99.999% purity, and the final firingcondition was 700° C. for 5 h. This powder was measured by the BETmethod to have a specific surface area of 45.4 m²/g, corresponding to anequivalent spherical particle diameter of 36.7 nm. Combustion analysisshowed that it had a residual carbon concentration of about 3 wt %. FIG.5 shows test results from electrodes and lithium half-cells constructedusing Swagelok hardware as in Example 1. It is seen that thefirst-charge capacity was lower than the first-discharge capacity by11.5%, both being measured at about a C/5 rate, showing that the initialnonstoichiometry of the sample may be about 11.5%. At higher C-rates,outstanding capacity retention was observed. At a 5C rate, the capacityretention was about 95%, at a 10C rate, the capacity retention was about90%, and at a 20C rate, the capacity retention was in the range 66-72%for three cells tested.

Example 3

Nanoscale ion storage materials having overall compositions LiFePO₄ andLi_(0.95)FePO₄ were synthesized and tested following procedures asdescribed in Example 2, with the mass of lithium carbonate beingadjusted so as to achieve the specified overall compositions. TheLiFePO₄ and Li_(0.95)FePO₄ powders were measured by the BET method tohave a specific surface areas of 39.78 m²/g and 46.2 m²/g respectively,corresponding to equivalent spherical particle diameters of 41.9 nm and36.1 nm respectively. Combustion analysis showed the two powders to bothhave residual carbon concentrations of 2.3 wt % and 3 wt % respectively.FIGS. 7 and 8 show the C/50 charge and discharge curves for these twosamples compared to a commercially purchased carbon-coated LiFePO₄ fromAldrich Chemical Company of several micrometer average particle size andmarkedly inferior rate capability. Due to the very high rate capabilityof these materials, see FIG. 19, these low-rate charge/discharge curvesshow the near-equilibrium voltages of the cells. From these curves it isseen that during continuous charge and discharge, a lithiumnonstochiometry x of at least about 15%, and y of at least about 10% isobtained. FIGS. 12-14 show PITT measurements of the nanoscaleLi_(0.95)FePO₄ sample as described earlier. During a single-stepdischarge to a voltage 5 mV above the OCV, 4.5% of the total dischargecapacity measured at C/50 rate of 160 mAh/g (3.8V to 2V) is measured,showing that greater than about 4.5% nonstoichiometry y can be obtainedduring dynamic discharging conditions. During a single-step charge to avoltage 5 mV below the OCV, 10.5% of the total charge capacity (2.9V to3.8V) was measured, showing that greater than about 10.5%nonstoichiometry x exists under dynamic charging conditions. Bycomparison, the y and x values measured from the capacities at 5 mVabove and below the OCV for the comparison sample from Aldrich Chemicalis only 0.7% and 1.2% respectively. FIG. 16 and Tables 1 and 2 showX-ray powder diffraction measurements of the nanoscale Li_(1-x)FePO₄sample as described earlier. From the Rietveld refinement of thissample, a crystallite size of about 28 nm was determined, which is closeto the calculated equivalent spherical particle size and shows that thehigh surface area of the sample is due to nanoscale crystallites of thelithium iron phosphate and not due to a high surface area impurity oradditive phase. FIG. 19 shows test results from three lithium half-cellsconstructed using Swagelok hardware as in Example 2.

Example 4

In this prophetic example, positive electrodes using a nanoscale ionstorage materials, for example, those of Examples 1 and 3 (having beenwell-characterized in their electrochemical performance over a widerange of C-rates), are used to construct a wound cylindrical lithium-ioncell. A high-rate graphite anode is employed, such as one utilizinggraphitized mesocarbon microbeads (MCMB, Osaka Gas Co.) of a fewmicrometers mean diameter. The performance of such cells, includingcharge capacity and energy at various C-rates, can be modeled from thevolumes and masses of the cell constituents when the density, thicknessand performance of individual electrodes in prototype cells is known, asin the present case. Starting from a fully charged state at 3.8V, anddischarging to a lower voltage limit of 2.0V, i.e., for 100% depth ofdischarge, the model shows that such cells will exhibit specific powerof at least about 500 W/kg (1000 W/L) at specific energy of at leastabout 100 Wh/kg (205 Wh/L), specific power of at least about 950 W/kg(2000 W/L) at specific energy of at least about 95 Wh/kg (190 Wh/L),specific power of at least about 1300 W/kg (2500 W/L) at specific energyof at least about 90 Wh/kg (180 Wh/L), and specific power of at leastabout 1600 W/kg (3200 W/L) at specific energy of at least about 85 Wh/kg(175 Wh/L). It is understood that for shallower depth of discharge, thespecific power and power density can be significantly higher than thesevalues.

As will be apparent to one of skill in the art from a reading of thisdisclosure, the present invention can be embodied in forms other thanthose specifically disclosed above. The particular embodiments describedabove are, therefore, to be considered as illustrative and notrestrictive. The scope of the invention is as set forth in the appendedclaims, rather than being limited to the examples contained in theforegoing description.

1. A lithium transition metal phosphate material for use as a positiveelectroactive material in a lithium ion battery having a specificsurface area of at least about 20 m²/g and comprising only onetransition metal, said lithium transition metal phosphate and saidsurface area selected to provide at least two co-existing olivine phasesduring cycling of the battery, wherein the two co-existing olivinephases include a lithium-rich transition metal phosphate phase and alithium-poor transition metal phosphate phase, wherein the percentagemolar volume difference between the two phases is less than about 6.4%.2. The lithium transition metal phosphate material of claim 1, whereinthe percentage molar volume difference between the two phases is lessthan about 6.25%.
 3. The lithium transition metal phosphate material ofclaim 1, wherein the percentage molar volume difference between the twophases is less than about 5.75%.
 4. The lithium transition metalphosphate material of claim 1, wherein the percentage molar volumedifference between the two phases is less than about 5.5%.
 5. Thelithium transition metal phosphate material of claim 1, wherein each ofthe two co-existing phases is crystalline and is defined by a unit cellhaving three lattice parameters, and wherein the differences in at leasttwo of the three lattice parameters between the two unit cells are lessthan 3%.
 6. The lithium transition metal phosphate material of claim 5,wherein the differences in all lattice parameters between the two unitcells are less than 4.7%.
 7. The lithium transition metal phosphatematerial of claim 5, wherein the differences in all lattice parametersbetween the two unit cells are less than 4.5%.
 8. The lithium transitionmetal phosphate material of claim 5, wherein differences in all latticeparameters between the two unit cells are less than 4.0%.
 9. The lithiumtransition metal phosphate material of claim 5, wherein the differencesin all lattice parameters between the two unit cells are less than 3.5%.10. The lithium transition metal phosphate material of claim 5, whereinthe two co-existing phases have a crystallographic plane in which themisfit strain between the two phases is less than about 1.6%.
 11. Thelithium transition metal phosphate material of claim 10, wherein themisfit strain between the two phases is less than about 1.5%.
 12. Thelithium transition metal phosphate material of claim 10, wherein themisfit strain between the two phases is less than about 1.4%.
 13. Thelithium transition metal phosphate material of claim 10, wherein none ofthe planes formed by any two of the lattice parameters of a unit cell ofthe crystal have a misfit strain between the two phases exceeding 8%.14. The lithium transition metal phosphate material of claim 13, whereinnone of the planes formed by any two of the lattice parameters of a unitcell of the crystal have a misfit strain between the two phasesexceeding 7.5%.
 15. The lithium transition metal phosphate material ofclaim 13, wherein none of the planes formed by any two of the latticeparameters of a unit cell of the crystal have a misfit strain betweenthe two phases exceeding 6.0%.
 16. The lithium transition metalphosphate material of claim 1, wherein the material has a specificsurface area of at least about 25 m²/g.
 17. The lithium transition metalphosphate material of claim 1, wherein the material has a specificsurface area of at least about 30 m²/g.
 18. The lithium transition metalphosphate material of claim 1, wherein the material has a specificsurface area of at least about 35 m²/g.
 19. The lithium transition metalphosphate material of claim 1, wherein the material has a specificsurface area of at least about 40 m²/g.
 20. The lithium transition metalphosphate material of claim 1, wherein the material has a specificsurface area of at least about 50 m²/g.
 21. The lithium transition metalphosphate material of claim 1, wherein material has an overallcomposition of Li_(1-x)MPO₄, where M comprises one first row transitionmetal selected from the group consisting of Ti, V, Cr, Mn, Fe, Co andNi, and wherein in use x ranges from 0 to
 1. 22. The lithium transitionmetal phosphate material of claim 21, wherein the composition is lithiumiron phosphate.
 23. The lithium transition metal phosphate material ofclaim 21, wherein the material exhibits a solid solution at roomtemperature over a composition range of 0<x<0.3, and/or 0.8<x<1.
 24. Thelithium transition metal phosphate material of claim 21, wherein thematerial exhibits a stable solid solution at room temperature over acomposition range of x between 0 and at least about 0.15 and/or 0.9<x<1.25. The lithium transition metal phosphate material of claim 21, whereinthe material exhibits a stable solid solution at room temperature over acomposition range of x between 0 and at least about 0.07 and/or0.95<x<1.
 26. The lithium transition metal phosphate material of claim21, wherein the lithium-rich transition metal phosphate phase has thecomposition Li_(y)MPO₄ exhibiting a stable solid solution at roomtemperature and the lithium-poor transition metal phosphate phase hasthe composition Li_(1-x)MPO₄ exhibiting a stable solid solution at roomtemperature, wherein 0.02<y<0.2 and 0.02<x<0.3.
 27. The lithiumtransition metal phosphate material of claim 26, wherein the solidsolution occupies a fraction of the compositional range of lithiumdefined as y+x.
 28. The lithium transition metal phosphate material ofclaim 1, wherein material has an overall composition ofLi_(1-x-z)M_(1-z)PO₄, where M comprises one first row transition metalselected from the group consisting of Ti, V, Cr, Mn, Fe, Co and Ni,where x is from 0 to 1 and z can be positive or negative.
 29. Thelithium transition metal phosphate material of claim 28, wherein thecomposition is lithium iron phosphate.
 30. The lithium transition metalphosphate material of claim 28, where z is between about 0.15 and −0.15.31. The lithium transition metal phosphate material of claim 1, whereinthe material is of a form selected from the group consisting ofparticles, agglomerated particles, fibers and coatings.
 32. The lithiumtransition metal phosphate material of claim 31, wherein the form has anaverage smallest cross-sectional dimension of about 75 nm or less. 33.The lithium transition metal phosphate material of claim 31, wherein theaverage smallest cross-sectional dimension is about 60 nm or less. 34.The lithium transition metal phosphate material of claim 31, wherein theaverage smallest cross-sectional dimension is about 45 nm or less. 35.The lithium transition metal phosphate material of claim 1, wherein thematerial is in the form of agglomerated particles and the averagecrystallite size as determined by x-ray diffraction is less than about800 nm.
 36. The lithium transition metal phosphate material of claim 35,wherein the average crystallite size as determined by x-ray diffractionis less than about 600 nm.
 37. The lithium transition metal phosphatematerial of claim 35, wherein the average crystallite size as determinedby x-ray diffraction is less than about 500 nm.
 38. The lithiumtransition metal phosphate material of claim 35, wherein the averagecrystallite size as determined by x-ray diffraction is less than about300 nm.
 39. The lithium transition metal phosphate material of claim 1,wherein the materials is crystalline.
 40. The lithium transition metalphosphate material of claim 1, wherein the material is amorphous.
 41. Acathode comprising the material of claim
 1. 42. An electrochemical cellcomprising the electrode of claim
 41. 43. A cathode comprising thematerial of claim
 21. 44. An electrochemical cell comprising theelectrode of claim
 43. 45. The lithium transition metal phosphatematerial of claim 1, wherein the material comprises disordered olivineformed upon delithiation or lithiation.
 46. The lithium transition metalphosphate material of claim 1, wherein the material comprises alithium-deficient lithium transition metal phosphate solid solutionformed upon delithiation at a temperature below 150° C.
 47. The lithiumtransition metal phosphate material of claim 46, wherein thelithium-deficient lithium transition metal phosphate solid solutioncomprises an ordered olivine structure.
 48. The lithium transition metalphosphate material of claim 47, wherein the deficiency occurs on thelithium or M1 sites of the ordered olivine.
 49. The lithium transitionmetal phosphate material of claim 47, wherein a disorder occurs on thelithium or M1 sites of the ordered olivine.
 50. The lithium transitionmetal phosphate material of claim 1, wherein the material transformsupon first charge to disordered olivine having a lithium deficient solidsolution and retains such solid solution at temperatures below 150° C.51. The lithium transition metal phosphate material of claim 50retaining the solid solution at temperatures below 100° C.
 52. Thelithium transition metal phosphate material of claim 50 retaining thesolid solution at temperatures below 50° C.